Note: Descriptions are shown in the official language in which they were submitted.
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High formability steel sheet for the manufacture of lightweight structural
parts and
manufacturing process
The invention relates to the manufacture of steel sheets or structural parts
combining a high elasticity modulus E in tension, a low density d and a high
processability, especially a high castability and high formability and
ductility.
The mechanical performance in stiffness of structural elements are known to
vary as
Ex/d, the coefficient x depending on the mode of external loading (for example
in tension
or in bending) and on the geometry of the elements (plates, bars). Thus,
steels exhibiting
both a high elasticity modulus and a low density have high mechanical
performances.
This requirement applies most particularly in the automotive industry, where
vehicle
lightening and safety are constant preoccupations. In order to produce steel
parts having
increased elasticity modulus and reduced density, it was proposed to
incorporate in the
steel ceramic particles of various types, such as carbides, nitrides, oxides
or borides.
Such materials indeed have a higher elasticity modulus, ranging from about 250
to 550
GPa, than that of base steels, which is around 210 GPa, into which they are
incorporated.
Hardening is achieved by load transfer between the steel matrix and the
ceramic particles
under the influence of a stress. This hardening is further increased due to
the matrix grain
size refinement by the ceramic particles. To manufacture these materials
comprising
ceramic particles distributed uniformly in a steel matrix, processes are known
that are
based on powder metallurgy: firstly, ceramic powders of controlled geometry
are
produced, these being blended with steel powders, thereby corresponding, for
the steel, to
an extrinsic addition of ceramic particles. The powder blend is compacted in a
mold and
then heated to a temperature such that this blend undergoes sintering. In a
variant of the
process, metal powders are blended so as to create the ceramic particles
during the
sintering phase.
This type of process however suffers from several limitations. Especially, it
requires
careful smelting and processing conditions in order not to cause a reaction
with the
atmosphere, taking into account the high specific surface area of metal
powders. Besides,
even after the compacting and sintering operations, residual porosities may
remain, such
porosities acting as damage initiation sites during cyclic stressing.
Furthermore, the
chemical composition of the matrix/particle interfaces, and therefore their
cohesion, is
difficult to control given the surface contamination of the powders before
sintering
(presence of oxides and carbon). In addition, when ceramic particles are added
in large
quantity, or when certain large particles are present, the elongation
properties decrease.
Finally, this type of process is suitable for low-volume production but cannot
meet the
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requirements of mass production in the automotive industry, and the
manufacturing costs
associated with this type of manufacturing process are high.
Manufacturing processes based on the extrinsic addition of ceramic powders
into the
liquid metal were also proposed. However, these processes suffer from most of
the
abovementioned drawbacks. More particularly, the difficulty of homogeneously
dispersing
the particles may be mentioned, such particles having a tendency to
agglomerate or to
settle in or float on the liquid metal.
Among the known ceramics that could be used to increase the properties of
steel is
in particular titanium diboride TiB2, which has the following intrinsic
characteristics:
Elasticity modulus: 583 GPa;
Relative density: 4.52.
In order to produce a steel sheet or part having increased elasticity modulus
and
reduced density, whilst avoiding the above mentioned problems, it was proposed
to
produce steel sheets having a composition with C, Ti and B contents such that
TiB2, Fe2B
and/or TiC precipitates form upon casting.
For example, EP 2 703 510 discloses a method for manufacturing a steel sheet
having a composition comprises 0.21% to 1.5% of 0,4% to 12% of Ti and 1.5% to
3% of
B, with 2.22*B Ti, the steel comprising TiC and TiB2 precipitates having an
average size
of below 10 pm . The steel sheets are produced by casting the steel in the
form of a semi-
product, for example an ingot, then reheating, hot-rolling and optionally cold-
rolling to
obtain a steel sheet. With such a process, elasticity modulus in tension
comprised
between 230 and 255 GPa can be obtained.
However, this solution also suffers from several limitations, arising both
from the
composition and from the manufacturing method, and leading to castability
issues, as well
as formability issues during the manufacturing process and during the
subsequent forming
steps performed on the steel sheet to produce a part:
- First, such steels have a low liquidus temperature (around 1300 C, so that
the
solidification starts at a relatively low temperature. In addition, the T1B2,
TiC and/or Fe2B
precipitate at an early stage of the casting process, at the beginning of the
solidification.
The presence of these precipitates and the low temperature result in a
hardening of the
steel and lead to rheological issues, not only during the casting process, but
also during
the further crop shearing and rolling operations. In particular, the
precipitates increase the
hot hardness of the solidified shell in contact with the mold, causing surface
defects and
increasing the risks of breakout. Consequently, surface defects, bleedings and
cracks
occur during the manufacturing process. In addition, owing to the high
hardness, the
range of achievable sizes for the hot-rolled or cold-rolled steel sheets is
limited. As an
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example, steel sheets 1 meter wide having a thickness lower than 3.5 mm cannot
be
produced in some hot strip mills due to rolling power limitation.
Second, despite the relatively small average size of the precipitates, the
size
distribution of the precipitates is wide. The steel thus comprises a
substantial fraction of
coarse precipitates, which negatively impact the formability, especially the
ductility and the
toughness of the steel, both during the manufacturing process of the sheet and
during the
subsequent forming operations to produce a part.
Besides, EP 1 897 963 discloses a method for manufacturing a steel sheet
having a
composition comprises 0.010% to 0.20% of C, 2.5% to 7.2% of Ti and 0.45xTi ¨
0.35%
B 0.45xTi +0.70%, the steel comprising TiB2 precipitates. However, this
document does
not address the problem of processability mentioned above.
Therefore, the invention aims at solving the above problems, in particular at
providing a steel sheet having an increased specific elasticity modulus in
tension together
with a high formability, especially a high ductility and a high toughness. The
invention also
aims at providing a manufacturing process of such a steel sheet, in which the
above
issues are not encountered.
The elasticity modulus in tension here designates the Young's modulus in the
transverse direction, measured by a dynamic Young's modulus measurement, for
example by a resonant frequency method.
The specific elasticity modulus in tension here refers to the ratio between
the
elasticity modulus in tension and the density of the steel. The density is for
example
determined using a helium pycnometer.
To that end, the invention relates to a steel sheet made of a steel having a
composition comprising, by weight percent:
0.010% C 0.080%
0.06% Mn 3%
Si 1.5%
0.005% 5 Al 1.5%
S 0.030%
P 0.040%,
Ti and B such that:
3.243/0 Ti 7.543/0
(0.45xTi) - 1.35 B (0.45xTi) - 0.43
optionally one or more elements chosen amongst:
Ni 1%
Mo 1%
4
Cr 3%
Nb 0.1%
V<_0.1%
the remainder being iron and unavoidable impurities resulting from the
smelting, said steel
sheet having a structure consisting of ferrite, at most 10% of austenite, and
precipitates, said
precipitates comprising eutectic precipitates of TiB2, the volume fraction of
TiB2 precipitates with
respect to the whole structure being of at least 9%, the proportion of TiB2
precipitates having a
surface area lower than 8 pm2 being of at least 96% said steel sheet
comprising no TiC
precipitates, or TiC precipitates with a volume fraction lower than 0.5%. the
steel sheet comprising
no Fe2B precipitates, the steel sheet having a content in free Ti of at least
0.95%.
Indeed, the inventors have found that with this composition, the content in
free Ti of the
steel is of at least 0.95%, and that owing to this content in free Ti, the
structure of the steel remains
mainly ferritic at any temperature below the liquidus temperature. As a
result, the hot hardness of
the steel is significantly reduced as compared to the steels of the state of
the art, so that the
castability and the hot formability are strongly increased.
In addition, the inventors have found that controlling the size distribution
of the TiB2
precipitates leads to a high formability, especially high ductility and
toughness, at high and low
temperatures, so that the hot and cold rollability of the steel is improved,
and parts with complex
shapes can be produced.
Preferably, the proportion of TiB2 precipitates having a surface area lower
than 3 pm2 is of at least
80%. Preferably, the proportion of TiB2 precipitates having a surface area
lower than 25 pm2 is of
100%.
Preferably, in the core region of the steel sheet, the proportion of TiB2
precipitates having a
surface area lower than 8 pm2 is of at least 96%, the proportion of TiB2
precipitates having a
surface area lower than 3 pm2 is preferably of at least 80% and the proportion
of TiB2 precipitates
having a surface area lower than 25 pm2 is preferably of 100%.
Preferably, the steel sheet comprises no TiC precipitates, or TiC precipitates
with a volume
fraction lower than 0.5% (with respect to the whole structure).
Generally, the steel sheet comprises no Fe2B precipitates.
According to an embodiment, the titanium, boron and manganese contents are
such that:
(0.45xTi) - 1.35 B (0.45xTi) ¨ (0.261*Mn) - 0.414.
According to an embodiment, the titanium and boron contents are such that:
(0.45xTi) - 1.35 B (0.45xTi) - 0.50.
According to an embodiment, the composition is such that C 0.050%.
Date Recue/Date Received 2021-05-26
4a
Preferably, the composition is such that Al 1.3%.
Date Recue/Date Received 2021-05-26
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Preferably, the steel sheet has a Charpy energy Kcv of at least 25 J/cm2 at -
40 C.
Generally, the steel sheet has a content in free Ti of at least 0.95%.
The invention also relates to a process for manufacturing a steel sheet, the
process
5 comprising the following successive steps:
- providing a steel having a composition comprising, by weight percent:
0.010% C 0.080%
0.06% Mn 3%
Si 1.5 /0
0.005% Al 1.5`)/0
S 0.030%
P 5 0.040%,
Ti and B such that:
3.2% Ti 7.5%
(0.45xTi) - 1.35 B (0.45xTi) - 0.43
optionally one or more elements chosen amongst:
Ni 1%
Mo 5 1%
Cr 3%
Nb 5 0 .1 cY0
V <0.1%
the remainder being iron and unavoidable impurities,
- casting the steel in the form of a semi-product, the casting temperature
being lower
than or equal to l_hquidus + 40 C, Lliquidus designating the liquidus
temperature of the steel,
the semi-product being cast in the form of a thin semi-product having a
thickness of at
most 110 mm, the steel being solidified during the casting with a
solidification rate
comprised between 0.03 cm/s and 5 cm/s at every location of the semi-product.
Indeed, the inventors have found that controlling cooling of the
solidification such
that the solidification rate is of at least 0.03 cm/s at every location of the
product,
especially at the core of the product, makes it possible to control the size
distribution of
the TiB2 precipitates. In addition, the casting under the form of a thin semi-
product, with
the composition of the invention allows achieving such high solidification
rates.
According to an embodiment, the semi-product is cast in the form of a thin
slab
having a thickness lower than or equal to 110 mm, preferably lower than or
equal to 70
mm.
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In an embodiment, the semi-product is cast in the form of a thin slab having a
thickness comprised between 15 mm and 110 mm, preferably between 15 mm and 70
mm, for example between 20 mm and 70 mm.
Preferably, the semi-product is cast by compact strip production.
According to another embodiment, the semi-product is cast in the form of a
thin strip
having a thickness lower than or equal to 6 mm, the solidification rate being
comprised
between 0.2 cm/s and 5 cm/s at every location of the semi-product.
Preferably, the semi-product is cast by direct strip casting between counter-
rotating
rolls.
Generally, after casting and solidification, the semi-product is hot rolled,
to obtain a
hot-rolled steel sheet.
Preferably, between casting and hot-rolling, the temperature of the semi-
product
remains higher than 700 C.
Preferably, before hot-rolling, the semi-product is de-scaled at a temperature
of at
least 1050 C.
According to an embodiment, after hot-rolling, the hot-rolled steel sheet is
cold-
rolled, to obtain a cold-rolled steel sheet having a thickness lower than or
equal to 2 mm.
Preferably, the titanium, boron and manganese contents are such that:
(0.45xTi) - 1.35 B (0.45xTi) ¨ (0.261*Mn) - 0.414.
Preferably, the composition is such that Al 1.3%.
The invention also relates to a method for manufacturing a structural part,
the
method comprising:
- cutting at least one blank from a steel sheet according to the invention or
produced
by a process according to the invention, and
- deforming said blank within a temperature range from 20 C to 900 C.
According to an embodiment, the method comprises, before deforming the blank,
a
step of welding the blank to another blank.
The invention also relates to a structural part comprising at least of portion
made of
a steel having a composition comprising, by weight percent:
0.010% C 0.080%
0.06% Mn 3%
Si 1.5 A
0.005`)/0 AI 1.5%
S 0.030%
P 0.040%,
Ti and B such that:
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3.2% Ti 7.5 /0
(0.45xTi) - 1.35 B (0.45xTi) - 0.43
optionally one or more elements chosen amongst:
Ni 1%
Mo 5 1%
Cr 3%
Nb 0.1%
V 0.1%
the remainder being iron and unavoidable impurities resulting from the
smelting,
said portion having a structure consisting of ferrite, at most 10% of
austenite, and
precipitates, said precipitates comprising eutectic precipitates of TiB2, the
volume fraction
of TiB2 precipitates with respect to the whole structure of said portion being
of at least 9%,
the proportion of TiB2 precipitates having a surface area lower than 8 pm2
being of at least
96%.
Preferably, the composition is such that Al 1.3%.
Preferably, the structural part is obtained by the method according to the
invention.
Other features and advantages of the invention will become apparent over the
course of the description below, given by way of non limiting example and with
reference
to the appended figures in which:
- Figure 1 is a micrograph illustrating the damage mechanism of individual
coarse
TiB2 precipitates,
- Figure 2 is a micrograph illustrating the damage mechanism of individual
fine 11B2
precipitates,
- Figure 3 is a micrograph illustrating fine TiB2 precipitates after a
collision of these
precipitates,
- Figure 4 is a micrograph illustrating coarse TiB2 precipitates after a
collision of
these precipitates,
- Figure 5 is a graph illustrating the reduction in area obtained through a
tensile test
at high temperatures for a steel of the invention and a comparative steel,
- Figure 6 is a micrograph illustrating the structure of a steel sheet
according to the
invention, along a longitudinal plane located at 1/4 of the thickness of the
steel sheet,
- Figures 7 and 8 are micrographs illustrating the structure of comparative
steel
sheets, along a longitudinal plane located at 1/4 of the thickness of the
steel sheets,
- Figure 9 is a micrograph illustrating the structure of the steel sheet of
Figure 6,
along a longitudinal plane located at half the thickness of the steel sheet,
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- Figures 10 and 11 are micrographs illustrating the structure of the
comparative
steel sheets of figures 7 and 8, along a longitudinal plane located at half
the thickness of
the steel sheets,
- Figure 12 illustrates the forming limit curves for the steel sheets of
Figures 6-11,
- Figures 13 and 14 are micrographs illustrating the damages of the steel
sheet of
Figures 7 and 10 after cold-rolling, along a longitudinal plane located at the
surface of the
cold-rolled steel sheet and along a longitudinal plane located at half the
thickness of the
cold-rolled steel sheet respectively,
- Figure 15 is a graph illustrating the Charpy energy Kcv of the steel
sheet of
Figures 6 and 9 and of the steel sheet of figures 8 and 11.
As regards the chemical composition of the steel, the carbon content is
adapted for
achieving the desired level of strength. For this reason, the carbon content
is of at least
0.010%.
However, the C content must be limited in order to avoid primary precipitation
of TiC
and/or Ti(C,N) in the liquid steel, and precipitation of TiC and/or Ti(C,N)
during eutectic
solidification and in the solid phase fraction, that could otherwise occur
owing to the high
Ti content of the steel. Indeed, TiC and Ti(C,N) precipitating in the liquid
steel would
deteriorate the castability by increasing the hot hardness of the solidified
shell during the
casting and lead to cracks in the cast product. In addition, the presence of
TiC precipitates
decreases the content in free Ti in the steel, and therefore inhibits the
alphageneous role
of Ti. For these reasons, the C content must be of at most 0.080%. Preferably,
the C
content is of at most 0.050%.
In a content of at least 0.06%, manganese increases the hardenability and
contributes to the solid-solution hardening and therefore increases the
tensile strength. It
combines with any sulfur present, thus reducing the risk of hot cracking.
However, if the
Mn content is higher than 3%, the structure of the steel will not be mainly
ferritic at all
temperatures, so that the hot hardness of the steel will be too high, as
explained in further
details below.
Silicon effectively contributes to increasing the tensile strength by solid
solution
hardening. However, excessive addition of Si causes the formation of adherent
oxides
that are difficult to remove by pickling, and the possible formation of
surface defects due in
particular to a lack of wettability in hot-dip galvanizing operations. To
ensure a good
coatability, the Si content must not exceed 1.5%.
In a content of at least 0.005%, aluminum is a very effective element for
deoxidizing
the steel. However, in a content above 1.5%, excessive primary precipitation
of alumina
occurs, impairing the castability of the steel.
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Preferably, the Al content is lower than or equal to 1.3%, so as to achieve a
further
improved castability.
In a content higher than 0.030%, sulfur tends to precipitate in excessively
large
amounts in the form of manganese sulfides, which reduce to a large extent the
hot and
cold formability of the steel. Therefore, the S content is of at most 0.030%.
Phosphorus is an element that segregates at the grain boundaries. Its content
must
not exceed 0.040% so as to maintain sufficient hot ductility, thereby avoiding
cracking,
and to prevent hot cracking during welding operations.
Optionally, nickel and/or molybdenum may be added, these elements increasing
the
tensile strength of the steel. For cost reasons, the additions of Ni and Mo
are each limited
to 1%.
Optionally, chromium may be added to increase the tensile strength, the Cr
content
being limited to at most 3% for cost reasons. Cr also promotes the
precipitation of borides.
However, the addition of Cr above 0.080% may promote the precipitation of (Fe,
Cr)
borides, to the detriment of TiB2 precipitates. Therefore, the Cr content is
preferably of at
most 0.080%.
Also optionally, niobium and vanadium may be added in an amount equal to or
less
than 0.1% so as to obtain complementary hardening in the form of fine
precipitated
carbonitrides.
Titanium and boron play an important role in the invention. Indeed, Ti and B
precipitate under the form of TiB2 precipitates which significantly increase
the elasticity
modulus in tension E of the steel. TiB2 may precipitate at an early stage of
the
manufacturing process, especially under the form of primary TiB2 precipitating
in the liquid
steel, and/or as eutectic precipitates.
However, the inventors have found that the TiB2 precipitates may lead to an
increase
in the hot hardness of the solidified shell during the casting and thereby
results in the
formation of cracks in the cast product, in the appearance of surface defects
and in a
decrease in the hot rollability of the steel which limit the accessible
thickness range for the
hot-rolled steel sheet.
Surprisingly, the inventors have found that if the Ti and the B content are
adjusted
such that the content of free Ti (hereinafter Ti*) is higher than or equal to
0.95%, the hot
hardness of the steel is significantly reduced. Indeed, the inventors have
found that under
this condition, the steel remains mainly ferritic, i.e. comprises at most 10%
of austenite,
whatever the temperature (below the liquidus), especially during the
solidification and the
hot-rolling, which leads to a decrease of the hot hardness of the steel as
compared to a
steel undergoing an allotropic transformation of more than 10% on cooling.
Thus, the
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castability and the hot ductility of the steel are improved to a large extent,
despite the
formation of TiB2 in the steel during solidification.
The "free Ti" here designates the content of Ti not bound under the form of
precipitates.
5 In
addition, a Ti* content of at least 0.95% greatly reduces, and even suppresses
the
formation of Fe2B that would impair the ductility.
Preferably, the Ti* content is higher than or equal to 0.92+0.58*Mn, wherein
Mn
designates the Mn content in the steel. Indeed, Mn is a gammageneous element
that may
favor the presence of austenite in the structure. Thus, the Ti* is preferably
adjusted
10
depending on the Mn content so as to ensure that the steel remains mainly
ferritic
whatever the temperature.
However, the Ti* content should remain lower than 3%, as no significant
beneficial
technical effect would be obtained from a Ti* content higher than 3%, despite
the higher
cost of adding titanium.
In order to ensure a sufficient TiB2 precipitation, and in the same time allow
the
content Ti* to reach 0.95%, the Ti content must be of at least 3.2%. If the Ti
content is
lower than 3.2%, the TiB2 precipitation is not sufficient, thereby precluding
a significant
increase in the elasticity modulus in tension, which remains lower than 220
GPa.
However, if the Ti content is higher than 7.5%, coarse primary TiB2
precipitation may
occur in the liquid steel and cause castability problems in the semi-product,
as well as a
reduction of the ductility of the steel leading to a poor hot and cold
rollability.
Therefore, the Ti content is comprised between 3.2% and 7.5%.
Besides, in order to ensure a Ti* content of at least 0.95%, the boron content
should
be of at most (0.45xTi) - 0.43, Ti designating the Ti content by weight
percent.
If B> (0.45xTi) - 0.43, the Ti* content will not reach 0.95%. Indeed, the Ti*
content
can be evaluated as Ti*= Ti ¨ 2.215x6, B designating the B content in the
steel. As a
consequence, if B > (0.45xTi) - 0.43, the structure of the steel will not be
mainly ferritic
during the casting and the hot rolling operations, so that its hot ductility
will be reduced,
which may lead to the formation of cracks and/or surface defects during the
casting and
hot rolling operations.
If a Ti* content higher than or equal to 0.92+0.58*Mn is targeted, the boron
content
should be of at most (0.45xTi) ¨ (0.261*Mn) - 0.414, Ti and Mn designating the
Ti and Mn
content by weight percent.
If B> (0.45xTi) ¨ (0.261*Mn) - 0.414, the Ti* content will not reach
0.92+0.58*Mn.
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The boron content should however be higher than or equal to (0.45xTi) - 1.35
to
ensure a sufficient precipitation of TiB2. In addition, a B content lower than
(0.45xTi) - 1.35
would corresponds to a Ti* content higher than 3%.
The balance is iron and residual elements resulting from the steelmaking.
According to the invention, the structure of the steel is mainly ferritic
whatever the
temperature (below Thquidus). BY "mainly ferritic", it must be understood that
the structure of
the steel consists of ferrite, precipitates (especially TiB2 precipitates) and
at most 10% of
austenite.
Thus, the steel sheet according to the invention has a structure which is
mainly
ferritic at all temperatures, especially at room temperature. The structure of
the steel sheet
at room temperature is generally ferritic, i.e. comprises no austenite.
The ferritic grain size is generally lower than 6 pm.
The volume fraction of TiB2 precipitates is of at least 9%, so as to obtain an
elasticity
modulus in tension E of at least 230 GPa.
The volume fraction of TiB2 precipitates is preferably of at least 12%, so as
to obtain
an elasticity modulus in tension E of at least 240 GPa.
The T1B2 precipitates mainly result from very fine eutectic precipitation upon
solidification, the mean surface area of the TiB2 precipitates being
preferably lower than
8.5 pm2, still preferably lower than 4.5 pm2, still preferably lower than 3
pm2.
The inventors have found that the size of the TiB2 precipitates in the steel
have an
influence on the properties of the steel, in particular on the damage
resistance of the
product during its manufacture, especially its hot and cold rollability, on
the damage
resistance of the steel sheet, especially during the forming operation, its
fatigue strength,
its fracture stress and its toughness.
However, the inventors have found that the main factor for ensuring a high
damage
resistance and therefore a high formability is the size distribution of the
TiB2 precipitates.
Indeed, the inventors have found that in a steel comprising TiB2 precipitates,
the
damages occurring during the manufacture, especially during the hot and/or
cold rolling
steps and the further forming operations, may result from damages undergone by
individual precipitates, and from collisions between the precipitates.
Especially, damage initiation of the individual T1B2 precipitates comes from
pile-up of
dislocations at the interface between the ferrite and the TiB2 precipitates,
and depends on
the size of the TiB2 precipitates. In particular, the fracture stress of the
TiB2 precipitates is
a decreasing function of the TiB2 precipitate size. If the size of some of the
TiB2
precipitates increases such that the fracture stress of these precipitates
becomes lower
than the interface disbanding stress, the damage mechanism changes from
interface
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disbonding to fracture of the TiB2 precipitates, leading to a significant
decrease of the
ductility, formability and toughness.
This change in damage mechanism is illustrated by Figures 1 and 2.
Figure 1 illustrates the damage of a coarse TiB2 precipitate under compressive
stress during cold-rolling: in that case, the TiB2 precipitate is fractured
along a direction
parallel to the compressive stress, under a relatively low stress.
By contrast, Figure 2 illustrates the interface disbonding of smaller TiB2
precipitates
during cold-rolling, by the appearance of cavities at the interface between
the ferritic
matrix and the TiB2 precipitates.
Consequently, if a steel sheet, though having TiB2 precipitates with a reduced
mean
size, comprises large TiB2 precipitates, these large TiB2 precipitates will
cause a change
in the damage mechanism of the steel and a decrease of the steel mechanical
properties.
Besides, the inventors have found that the damages resulting from collisions
between TiB2 precipitates are all the more important that the size of these
precipitates is
large. In particular, whereas a collision between coarse TiB2 precipitates
results in a
fracture of these precipitates, a collision of small TiB2 precipitates does
not lead to such
fracture.
Figures 3 and 4 illustrate precipitates of different sizes further to a
collision.
Especially, Figures 3 and 4 illustrate fine precipitates and large TiB2
precipitates
after a collision respectively. These figures show that the collision of the
large precipitates
led to a fracture of one of the colliding precipitates, whereas the collision
of the fine
precipitates did not result in any damage.
In order to ensure high ductility, formability and toughness, the inventors
have found
that the distribution of the size of the TiB2 precipitates must be such that
the proportion of
TiB2 precipitates having a surface area lower than 8 m2 is of at least 96%.
Moreover, the proportion of TiB2 precipitates having a surface area lower than
3 pm2
should preferably be of at least 80%, and the proportion of TiB2 precipitates
having a
surface area lower than 25 m2 should preferably be of 100%.
The proportion of TiB2 precipitates having a surface area lower than 3 pm2, 8
pm2 or
25 m2 is defined as the number of TiB2 precipitates having a surface area
lower than 3
m2, 8 m2 or 25 pm2, divided by the number of T1B2 precipitates, and
multiplied by a
factor 100.
The proportion of TiB2 precipitates having a surface area lower than 3 pm2, 8
pm2 or
25 m2 is preferably determined on a specimen prepared using standard
metallographic
technique for surface preparation and etched with nital reagent, by image
analysis using a
Scanning Electron Microscope (S EM).
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Especially, at the core of the sheet, the distribution of the size of the TiB2
precipitates
must be such that the proportion of TiB2 precipitates having a surface area
lower than 8
pm2 is of at least 96%, and preferably such that the proportion of TiB2
precipitates having
a surface area lower than 3 pm2 is of at least 80%, still preferably such that
the proportion
of TiB2 precipitates having a surface area lower than 25 pm2 is of 100%.
By considering a sheet having a generally rectangular shape having a length //
in a
longitudinal direction, a width w1 in the transversal direction and a
thickness t/ in the
thickness direction, the core of the sheet is defined as the portion of the
sheet extending
over the length // and over the width w/, in the thickness direction of the
sheet, from a
first end located at 45% of the overall thickness t1 of the sheet to a second
end located at
55% of the overall thickness t1 of the sheet.
Indeed, the inventors have found that under this condition, the damages occur
by
interface disbonding, so that the damage kinetics is delayed. Besides, under
this
condition, the damages that may result from collisions between T1B2
precipitates are
highly reduced.
As a consequence, the formability and the ductility of the steel sheet during
its
manufacture and in use are greatly improved.
In particular, the reduction ratio achievable through cold-rolling is
increased, and the
formability is increased, so that parts with complex shapes can be formed.
Having a proportion of TiB2 precipitates having a surface area lower than 8
pm2 of at
least 96% is critical. Indeed, the inventors have found that below this value,
the coarse
TiB2 precipitates cause a change in damage mechanism, as explained above,
which
drastically reduces the damage resistance of the steel.
Besides, the steel sheet according to the invention comprises no or a small
fraction
of TiC precipitates, the volume fraction of TiC precipitates in the structure
remaining lower
than 0.5%, generally lower than 0.36%.
Indeed, as explained above, TiC precipitates, if present, would have formed in
the
liquid steel, and would have deteriorated the castability of the steel, so
that a fraction of
TiC precipitates in the structure higher than 0.5% would result in cracks
and/or surface
defects in the steel sheet. The presence of TiC precipitates further decreases
the ductility
of the steel.
In addition, owing to the high Ti* content, the steel sheet does not comprise
any
Fe2B precipitates, the volume fraction of Fe2B precipitates in the structure
being of 0%.
The absence of Fe2B precipitates increases the ductility of the steel sheet.
The steel sheet, whether hot-rolled or cold-rolled, has a very high toughness,
even
at low temperatures. Especially, the transition temperature from ductile mode
to mixed
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mode is lower than -20 C, and the Charpy energy Kcv of the steel sheet is
generally
higher than or equal to 25 J/cm2at -40 C, and higher than or equal to 20
J/cm2at -60 C.
The steel sheet has an elasticity modulus in tension E of at least 230 GPa,
generally
of at least 240 GPa, a tensile strength TS of at least 640 MPa and a yield
strength of at
least 250 MPa before any skin-pass. Thus, a non skin-passed sheet according to
the
invention generally has a yield strength of at least 250 MPa.
The high tensile strength, of at least 640 MPa, is especially achieved owing
to the
small size and the size distribution of the TiB2 precipitates in the steel of
the invention, due
to the Hall-Petch effect and increased work-hardening.
The elasticity modulus in tension is an increasing function of the fraction of
TiB2
precipitates.
Especially, an elasticity modulus in tension E of at least 230 GPa is achieved
with a
fraction of TiB2 precipitates of 9% or higher. In the preferred embodiment
wherein the
volume fraction of T1B2 precipitates is of at least 12%, an elasticity modulus
in tension E of
at least 240 GPa is achieved.
Besides, the presence of TiB2 precipitates leads to a decrease of the density
of the
steel.
As a consequence, the steel sheet of the invention has a very high specific
elasticity
modulus in tension.
A process for manufacturing a steel sheet according to the invention is
implemented
as follows.
A steel with the composition according to the invention is provided, and the
steel is
then cast into a semi-product.
The casting is performed at a temperature lower than or equal to Tliquidus+40
C,
Thquidus designating the liquidus temperature of the steel.
Indeed, a casting temperature higher than Thquidus + 40 C could lead to the
formation
of coarse TiB2 precipitates.
The liquidus temperature Tliquidus of the steel of the invention is generally
comprised
between 1290 C and 1310 C. Therefore, the casting temperature should generally
be of
at most 1350 C.
The casting is carried out so as to form upon casting a thin product, having a
thickness of at most 110 mm, especially a thin slab or a thin strip.
To that end, the casting is preferably performed by compact strip production,
to form
a thin slab having a thickness lower than or equal to 110 mm, preferably of at
most 70
mm, or by direct strip casting between counter-rotating rolls, to form a thin
strip having a
thickness lower than or equal to 6 mm.
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In any case, the thickness of the semi-product must be of at most 110 mm, and
preferably of at most 70 mm.
For example, the semi-product is cast in the form of a thin slab having a
thickness
comprised between 15 mm and 110 mm, preferably between 15 mm and 70 mm, for
5 example between 20 mm and 70 mm.
Casting the semi-product under the form of a thin semi-product, for example a
thin
slab or strip, improves the processability of the steel by limiting the damage
of the steel
during rolling and forming operations.
Indeed, casting the semi-product under the form of a thin semi-product, for
example
10 a thin slab or strip allows using during the subsequent rolling steps a
lower reduction rate
to achieve the desired thickness.
A decrease in the reduction rate limits the damage of the steel that may
result from
collisions of the TiB2 precipitates during hot and cold rolling operations.
Most of all, the casting under the form of a thin semi-product allows
achieving very
15 fine TiB2 precipitates, so that the damage that may result from
collisions of TiB2
precipitates and the damage of individual TiB2 precipitates are reduced, as
explained
above.
Especially, the casting under the form of a thin semi-product allows a fine
control of
the solidification rate upon cooling across the thickness of the sheet,
ensures a
solidification rate fast enough in the whole product and minimizes the
difference in
solidification rate between the surface of the product and the core of the
product.
Indeed, achieving a sufficient and homogeneous solidification rate is
necessary to
obtaining very fine TiB2 precipitates, not only at the surface of the product,
but also at the
core of the semi-product. By considering a semi-product having a generally
rectangular
shape having a length 12 in a longitudinal direction, a width w2 in the
transversal direction
and a thickness t2 in the thickness direction, the core (or core region) of
the semi-product
is defined as the portion of the semi-product extending over the length 12 and
over the
width w2, in the thickness direction of the semi-product, from a first end
located at 45% of
the overall thickness t2 of the semi-product, to a second end located at 55%
of the overall
thickness of the semi-product.
The inventors have further found that in order to obtain very fine TiB2
precipitates
such that the proportion of TiB2 precipitates having a surface area lower than
8 pm2 is of
at least 96%, the cooling conditions during the solidification must be such
that the steel is
solidified with a solidification rate equal to or greater than 0.03 cm/s, up
to 5 cm/s, at every
location of the semi-product.
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Owing to the decrease of the solidification rate from the surface to the core
of the
product, a solidification rate of at least 0.03 cm/s at every location implies
that the
solidification rate at the core of the product is of at least 0.03 cm/s, up to
5 cm/s.
Besides, if the semi-product is cast under the form of a thin strip,
especially by direct
strip casting between counter-rotating rolls, to form a thin strip having a
thickness lower
than or equal to 6 mm, the solidification rate is comprised between 0.2 cm/s
and 5 cm/s at
every location of the semi-product.
Indeed, the inventors have found that a solidification rate of at least 0.03
cm/s at
every location, especially at the core of the product, allows obtaining very
fine TiB2
precipitates, not only at the surface of the product but also throughout the
whole thickness
of the product, such that the mean area surface is lower than 8.5 pm2 and the
proportion
of TiB2 precipitates having a surface area lower than 8 pm2 is of at least
96%. In addition,
the proportion of TiB2 precipitates having a surface area lower than 3 pm2 is
of at least
80%, and the proportion of TiB2 precipitates having a surface area lower than
25 pm2 is of
100%.
Especially, a solidification rate of at least 0.03 cm/s in the core region of
the product
allows obtaining very fine TiB2 precipitates in the core region of the semi-
product, such
that the mean area surface is lower than 8.5 pm2 and the proportion of TiB2
precipitates
having a surface area lower than 8 pm2 is of at least 96%. In addition, the
proportion of
T1B2 precipitates having a surface area lower than 3 pm2 is of at least 80%,
and the
proportion of TiB2 precipitates having a surface area lower than 25 pm2 is of
100%.
By contrast, if the solidification rate at least some parts of the product is
lower than
0.03 cm/s, TIC precipitates and/or coarse TiB2 precipitates will form during
solidification.
The control of the cooling and solidification rates to the above values is
achieved
owing to the casting of the steel in the form of a thin semi-product with a
thickness lower
than 110 mm, and to the composition of the steel.
Especially, the casting in the form of a thin semi-product results in a high
cooling
rate across the product thickness and in an improved homogeneity of the
solidification
rate from the surface to the core of the product.
In addition, owing to the high Ti* content of the steel, the steel solidifies
mainly as
ferrite. Especially, the solidified steel has a mainly ferritic structure from
the start of
solidification and during the whole solidification process, the austenite
fraction in the steel
remaining of at most 10%. Thus, no or very limited phase transformation occurs
during the
cooling.
As a result the steel can be cooled by rewetting, rather than by film boiling,
which
allows reaching very high solidification rates.
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Film boiling is a cooling mode in which a thin layer of vapor of cooling
fluid, having a
low thermal conductivity, is interposed between the surface of the steel and
the liquid
cooling fluid. In film boiling, the heat transfer coefficient is low. By
contrast, cooling by
rewetting occurs when the vapor layer is fractured, and the cooling fluid
becomes in
contact with the steel. This cooling mode occurs when the temperature of the
surface of
the steel is lower than the Leidenfrost temperature. The heat transfer
coefficient achieved
through rewetting is higher than the heat transfer coefficient achievable
through film
boiling, so that the solidification rate is increased. However, if phase
transformations occur
during cooling by rewetting, the coupling between rewetting and phase
transformation
induces high strains in the steel resulting in cracks and surface defects.
Therefore, steels enduring a significant allotropic transformation during
solidification
cannot be cooled by rewetting.
By contrast, in the steels of the invention, which comprise at most 10% of
austenite
at any temperature, little or no phase transformation occurs upon
solidification, and the
steel can therefore be cooled by rewetting.
Thus, very high solidification rates can be achieved.
At the end of the solidification, the structure of the steel is mainly
ferritic and
comprises very fine eutectic Ti B2 precipitates.
In addition, owing to the mainly ferritic structure of the steel as soon as
the
solidification starts, no or little transformation of 8 ferrite into austenite
occurs during
solidification (i.e. at most 10% of 8 ferrite transforms into austenite during
solidification), so
that the local contractions that would result from this transformation, which
could lead to
cracks in the semi-product, are avoided.
In particular, in the absence of significant transformation of 6 ferrite into
austenite, no
peritectic induced precipitation occurs during solidification. Such peritectic
induced
precipitation, occuring in the dendrites, could lead to a decrease of the hot
ductility and
induce cracks, especially during the further hot rolling.
Therefore, the solidified semi-product has a very good surface quality and
comprises
no or very few cracks.
Moreover, the solidification of the steel as mainly ferrite, as compared to a
structure
comprising more than 10% of austenite at the solidification, reduces to a
large extent the
hardness of the solidified steel, in particular the hardness of the solidified
shell.
Especially, the hardness of the steel is about 40% lower than a comparable
steel
that would have an structure comprising more than 10% of austenite during
solidification.
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The low hot hardness of the solidified steel results in a reduction of the
rheological
issues involving the solidified shell, especially avoids the occurrence of
surface defects,
depression and bleedings in the cast product.
In addition, the low hot hardness of the solidified steel also guarantees a
high hot
ductility of the steel, as compared to allotropic grades.
Owing to the high hot ductility of the product, the formation of cracks, that
would
otherwise appear during the bending and straightening operations of the
casting process,
and/or during the subsequent hot rolling, is avoided.
After solidification, the semi-product is cooled to an end of cooling
temperature
which is preferably of not less than 700 C. At the end of the cooling, the
structure of the
semi-product remains mainly ferritic.
The semi-product is then heated, from the end of cooling temperature to about
1200 C, de-scaled then hot-rolled.
During de-scaling, the temperature of the surface of the steel is preferably
of at least
1050 C. Indeed, below 1050 C, liquid oxides will solidify on the surface of
the semi-
product, which may cause surface defects.
Preferably, the semi-product is directly hot-rolled, i.e. is not cooled to a
temperature
below 700 C before hot-rolling, such that the temperature of the semi-product
remains at
any time higher than or equal to 700 C between the casting and the hot-
rolling. The direct
hot-rolling of the semi-product allows reducing the time necessary for
homogenizing the
temperature of the semi-product before hot-rolling, and therefore limiting the
formation of
liquid oxides at the surface of the semi-product.
In addition, the as cast semi-product is generally brittle at low
temperatures, so that
directly hot-rolling the semi-product allows avoiding cracks that may
otherwise occur at
low temperatures due to the brittleness of the as cast semi-product.
The hot-rolling is for example performed in a temperature range comprised
between
1100 C and 900 C, preferably between 1050 C and 900 C.
As explained above, the hot ductility of the semi-product is very high, owing
to the
mainly ferritic structure of the steel. Indeed, no or little phase
transformation, which would
reduce the ductility, occurs in the steel during hot-rolling.
As a consequence, the hot rollability of the semi-product is satisfactory,
even with a
hot-rolling finish temperature of 900 C, and the appearance of cracks in the
steel sheet
during hot-rolling is avoided.
For example, hot-rolled steel sheets having a thickness comprised between 1.5
mm
and 4 mm, for example comprised between 1.5 mm and 2 mm, are obtained.
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After hot-rolling, the steel sheet is preferably coiled. The hot-rolled steel
sheet is
then preferably pickled, for example in an HCI bath, to guarantee a good
surface quality
Optionally, if a lower thickness is desired, the hot-rolled steel sheet is
subjected to
cold-rolling, so as to obtain a cold-rolled steel sheet having a thickness of
less than 2 mm,
for example comprised between 0.9 mm and 1.2 mm.
Such thicknesses are achieved without producing any significant internal
damage.
This absence of significant damage is especially due to the casting under the
form of a
thin semi-product and to the composition of the steel.
Indeed, since the cold-rolled sheet is produced from a thin product, the hot
and cold
reduction ratios necessary to achieve a given thickness is reduced. Therefore,
the
occurrence of collisions between the T1B2 precipitates, which could lead to
damage, is
reduced.
Furthermore, owing to the size distribution of the TiB2 precipitates, achieved
thanks
to the low thickness of the semi-product and to the composition, cold
reduction ratios of up
to 40%, and even of up to 50% can be achieved without producing any
significant internal
damage.
Indeed, since the steel comprises no coarse TiB2 precipitates, the damages
occur by
interface disbonding, so that the damage kinetics is delayed. Besides, the
collision of the
TiB2 precipitates, owing to their small sizes, does not lead to any
significant damage.
As a consequence, the occurrence of damages during cold-rolling is highly
reduced.
After cold-rolling, the cold-rolled steel sheet may be subjected to an
annealing. The
annealing is for example performed by heating the cold-rolled steel sheet at a
mean
heating rate preferably comprised between 2 and 4 C/s, to an annealing
temperature
comprised between 800 C and 900 C, and holding the cold-rolled steel sheet at
this
annealing temperature for an annealing time generally comprised between 45 s
and 90s.
The steel sheet thus obtained, which may be hot-rolled or cold rolled, has a
mainly
ferritic structure, i.e. consists of ferrite, at most 10% of austenite, and
precipitates.
Generally, the steel sheet thus obtained has a ferritic structure at room
temperature, i.e. a
structure consisting of ferrite and precipitates, without austenite.
The steel sheet thus obtained comprises TiB2 precipitates, which are eutectic
TiB2
precipitates, the volume fraction of T1B2 precipitates being of at least 9%.
The proportion of TiB2 precipitates in the steel sheet having a surface area
lower
than 8 m2 is of at least 96%. In addition, the proportion of TiB2
precipitates having a
surface area lower than 3 m2 is preferably of at least 80%, and the
proportion of TiB2
precipitates having a surface area lower than 25 m2 is preferably of 100%.
This is especially the case in the core region of the sheet.
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The steel sheet thus obtained comprises a very small amount of TiC
precipitates,
owing to the low C content of the steel and to the manufacturing process, and
to the
absence of peritectic induced precipitation during solidification. The volume
fraction of TIC
precipitates in the structure is in particular lower than 0.5%, generally
lower than 0.36%.
5 The steel sheet thus obtained comprises no Fe2B precipitates.
With this manufacturing process, the formation of surface defects and cracks
in the
cast product and the steel sheet is avoided.
Especially, the reduction in hardness achieved owing to the high Ti* content
allows
avoiding the occurrence of surface defects, depression and bleedings in the
cast product.
10 In addition, the steel sheet thus obtained has very high formability,
toughness and
fatigue strength, so that the parts with complex geometry can be produced from
such
sheets.
Especially, the damages in the steel sheet that may result from hot and/or
cold-
rolling are minimized, so that steel has an improved ductility during the
subsequent
15 forming operations and an improved toughness.
Furthermore, the high elasticity modulus in tension of the steel according to
the
invention reduces the springback after the forming operations and thereby
increases the
dimensional precision on the finished parts.
To produce a part, the steel sheet is cut to produce a blank, and the blank is
20 deformed, for example by drawing or bending, in a temperature range
comprised between
20 and 900 C.
Advantageously, structural elements are manufactured by welding a steel sheet
or
blank according to the invention to another steel sheet or blank, having an
identical or a
different composition, and having an identical or a different thickness, so as
to obtain a
welded assembly with varying mechanical properties, which can be further
deformed to
produce a part.
For example, the steel sheet according to the invention may be welded to a
steel
sheet made of a steel having a composition comprising, by weight percent:
0.01% C 5 0.25%
0.05% Mn 2%
Si 5 0.4%
Al 5 0.1%
Ti 0.1%
Nb 0.1%
V 0.1%
Cr 3%
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MO 5 1%
Ni 1%
B 5 0.003%
the remainder being iron and unavoidable impurities resulting from the
smelting.
Examples:
As examples and comparison, sheets made of steel compositions according to
table
I, have been manufactured, the elements being expressed in weight percent.
Ti*=Ti-
Mn Si Al S P Ti B Cr
2.215*B
A 0.0227 0.061 0.168 0.039 0.0067 0.008 5.32 1.67 0.12 1.6
B 0.04 0.09
0.14 0.146 0.0015 0.009 6.34 2.34 0.075 1.16
C 0.036 0.07 0.15 0.065 0.001 0.01 5.3 2.05 0.05 0.75
Table 1
In Table 1, the underline value is not according to the invention.
These steels were cast in the form of semi-products:
- steel A was continuously cast in the form of a slab having a thickness of
65 mm
(sample 11),
- steel B was cast in the form of a ingot of 300 kg, having a section of 130
mm x 130
mm (sample R1),
- steel C was cast in the form of a thin slab having a thickness of 45 mm
(sample
R2).
The solidification rates during solidification of the cast products were
assessed at the
surface and at the core of the products, and are reported in Table 2 below.
Sample Steel At the surface At the core
reference composition (cm/s) (cm/s)
11 A 0.3 0.06
R1 B 0.001 0.0001
R2 C 0.3 0.01
Table 2
In Table 2, the underlined values are not according to the invention.
Sample 11 was cast under the form of a thin slab, having a thickness lower
than 110
mm.
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In addition, the composition (A) of sample 11 is in accordance with the
invention, and
has therefore a content in free Ti of at least 0.95%, so that during the
solidification, no or
little phase transformation occurred, allowing cooling by rewetting.
Owing to the low thickness of the cast product and to the cooling by
rewetting, the
solidification rate for sample 11 could be higher than 0.03 cm/s, even at the
core of the
semi-product.
By contrast, sample R1 has a composition (B) according to the invention, but
was
not cast as a thin semi-product, its thickness being higher than 110 mm.
As a consequence, the solidification rate could not reach the targeted values,
neither
at the core nor at the surface of the semi-product.
Sample R2 does not have a composition (C) in accordance with the invention,
its B
content being higher than (0.45xTi) - 0.43. Thus, sample R2 has a content in
free Ti lower
than 0.95% (0.75%).
Thus, even if the steel was cast under the form of a thin strip, an important
phase
transformation occurring during solidification, so that the cooling could not
be performed
by rewetting. As a result, the solidification rate did not reach 0.03 cm/s at
the core of the
product.
The inventors have investigated the hot formability of samples 11 and R2.
Especially, the hot formability of as cast samples 11 and R2 was assessed by
performing hot plane strain compression tests with various strain rates as
temperatures
ranging from 950 C to 1200 C.
To that end, Rastegaiev specimens were sampled from as cast samples 11 and R2.
The specimens were heated to a temperature of 950 C, 1000 C, 1100 C or 1200 C,
and
then compressed by two punches, located of opposite sides of the specimen,
with various
strain rates of 0.1 s-1, 1 s-1, 10 s-1 or 50 s-1. The stresses were
determined, and for each
test, the maximum stress was assessed.
Table 3 below reports at each temperature and for each of the samples 11 and
R2
the fraction of austenite in the structure at this temperature, and the
maximal stress
determined at each temperature for each strain rate.
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950 C 1000 C 1100 C 1200 C
11 R2 11 R2 11 R2 11 R2
% of
<10% 100% <10% 100% <10% 100% <10% 100%
austenite
Strain
Stress max (MPa)
rate (s-1)
0.1 93 196 70 169.5 47 127 27 81
1 138 230 108 209 75 164 53 112
199 270 169 253 125 212 90 153
50 236 316 204 294 155 250 126 191
Table 3
These results show that the maximum stress reached for sample 11 is much lower
than for sample R2, whatever the temperature comprised between 950 C and 1200
C and
whatever the strain rate, the maximum stress for steel 11 being of up to 67%
lower than
5 the maximum stress reached for steel R2.
This reduction of the maximum stress results especially from the difference
between
the structure of sample 11, which is mainly ferritic at all temperatures, and
the structure of
sample R2, which endures phase transformation and becomes austenitic at high
temperatures. This reduction implies that at high temperatures, the hardness
of the steel
10 of the invention is reduced to a large extent as compared to a steel
having a Ti* content
lower than 0.95%, the hot formability being thereby improved.
The hot formability of as cast samples 11 and R2 was further assessed by
performing high temperature tensile test on a thermomechanical simulator
Gleeble.
Especially, the reduction of area was determined at temperatures ranging from
600 C to 1100 C.
The results of these tests, illustrated on Figure 5, show that the hot
ductility of
sample 11 remains high even at decreasing temperatures, especially at
temperatures
comprised between 800 C and 900 C, whereas the ductility of sample R2
drastically
decreases with the temperature.
As a consequence, sample 11 can be processed at lower temperatures than sample
R2. Conversely, during the manufacturing process, the occurrence of cracks or
bleedings
in sample 11 will be largely reduced as compared to sample R2.
The inventors have further characterized the TiB2 precipitates of the as cast
products
on samples taken from I/4 the thickness from samples 11, R1 and R2, and a
sample taken
from half the thickness of sample 11 by image analysis using a Scanning
Electron
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Microscope (SEM). The specimens for microscopic examination were prepared
using
standard metallographic technique for surface preparation and etched with
nital reagent.
The size distributions are reported in Table 4 below
As shown in Table 4, sample R1 comprise a high percentage of coarse
precipitates,
having a surface area higher than 8 pm2.
Sample R2 comprises a higher fraction of small TiB2 precipitates than sample
R1.
However, the percentage of TiB2 precipitates having a surface area lower than
8 pm2 for
sample R2 does not reach 96%.
By contrast, sample 11 has a very high fraction of TiB2 precipitates with an
area of at
most 8 pm2, especially higher than 96% In addition, the fraction of TiB2
precipitates with
an area of at most 3 pm2 is higher than 80%, and all the TiB2 precipitates
have an area
lower than or equal to 25 pm2.
Percentage of Percentage of Percentage of
Sample TiB2 with an TiB2 with an TiB2 with an
reference area of at most area of at most area
of at most
3 prre 8 pm2 25 pm2
11 83.9 96.7 100
R1 46.6 70.3 86.7
R2 81.2 94.5 98.5
Table 4
In Table 4, the underlined values are not according to the invention.
Besides, after solidification, sample 11 was heated to a temperature of 1200
C, then
hot-rolled with a final rolling temperature of 920 C, to produce a hot-rolled
sheet having a
thickness of 2.4 mm.
The hot-rolled steel sheet 11 was further cold-rolled with a reduction ratio
of 40% to
obtain a cold-rolled sheet having a thickness of 1.4 mm.
After cold-rolling, the steel sheet 11 was heated with an average heating rate
of
3 C/s to an annealing temperature of 800 C and held at this temperature for 60
s.
After solidification, samples R1 and R2 were cooled to room temperature, then
reheated to a temperature of 1150 C and hot-rolled with a final rolling
temperature of
920 C to produce a hot-rolled sheet having a thickness of 2.2 mm and 2.8 mm
respectively.
The microstructures of the hot-rolled sheets produced from samples 11, R1 and
R2
were investigated by collecting samples at locations situated at 1/4 the
thickness of the
sheets and at half the thickness of the sheets, so as to observe the structure
along
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longitudinal plane at half distance between the core and the surface of the
sheets and at
the core of the sheets respectively.
The microstructures were observed with a Scanning Electron Microscope (SEM)
after etching with the Klemm reagent.
5 The microstructure of steels 11, R1 and R2 at 1/4 of the thickness are
shown on
Figures 6, 7 and 8 respectively.
The microstructure of steel sheets 11, R1 and R2 at half the thickness are
shown on
Figures 9, 10 and 11 respectively.
These figures show that the structure of steel 11 is very fine, both at 1/4
thickness and
10 at the core of the product.
By contrast the structure of steel R1, which was cooled with lower
solidification
rates, comprises coarse grains.
The structure of steel R2, though comprising fine grains at 1/4 thickness,
also
comprises coarse grains, especially at the core of the semi-product.
15 Overall, the structure of steel 11 is very homogeneous, whereas the
structures of
steels R1 and R2 each comprise grains with very different sizes.
The inventors have further investigated the cold formability of steels 11, R1
and R2.
The cold formability of the steels was assessed on steels sheets produced from
as
cast steels 11, R1 and R2 with plane strain tests.
20 Especially, samples were collected from the sheets made of steels 11, R1
and R2,
and the forming limit curves for steels 11, R1 and R2 were determined. These
forming limit
curves are illustrated on Figure 12, and the measurements reported in Table 5
below.
As shown by Figure 12 and Table 5, steel 11 has an improved formability as
compared to steels R1 and R2.
25 Without being bound to a theory, it is thought that the presence of
coarse TiB2
precipitates in steels R1 and R2, even in a small quantity, promotes
localization of the
strain during the forming operations, in the present case during the bending,
which leads
to a poorer formability than steel 11. It is further thought that the
localization may result
from the early damage of the coarse TiB2 precipitates colliding.
By contrast, steel 11 comprise no coarse precipitates, which minimizes the
collision
of the TiB2 precipitates and therefore improves the formability.
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26
Steel
E2 El
-0.061 0.292
-0.052 0.275
0.007 0.224
0.02 0.229
11 0.031 0.2
0.034 0.247
0.047 0.205
0.058 0.212
0.062 0.24
0.00718 0.165
R1 0.00821 0.161
0.0103 0.136
0.016 0.104
R2 0.017 0.107
0.021 0.111
0.023 0.144
Table 5
To confirm the influence of the size of the TiB2 precipitates on the
formability, the
inventors subjected a hot-rolled steel sheet R1, obtained through the process
disclosed
above, to cold-rolling, with a cold reduction ratio of 50%. After cold-
rolling, the steel sheet
R1 was heated with an average heating rate of 3 C/s to an annealing
temperature of
800 C and held at this temperature for 60 s.
The inventors then collected specimens from the surface and from the core of
the
cold-rolled steel sheet R1 (after annealing), and observed these specimens by
Scanning
Electron Microscopy.
The structures observed at the surface and at the core are illustrated on
Figures 13
and 14 respectively.
As visible on these figures, the specimen collected from the surface of the
sheet
comprises few damages, unlike the specimen collected form the core, in which
an
important damaging is observed.
These observations confirm that the coarse TiB2 precipitates, which are mainly
located at the core of the sheet owing to the lower solidification rate in
this portion, cause
damage during deformation and therefore degrade the formability of the steel.
The bending ability of steels 11, R1 and R2 was assessed by performing an edge
bending test (also named 90 flanging test) on samples collected from the hot-
rolled steel
sheets made of steels Ii, Al and R2, and from the cold-rolled steel sheet
(after annealing)
made of steel 11.
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The samples were held between a pressure pad and a die, and a sliding die was
slid to bend the portion of the sample protruding from the pad and the die.
The bending
test was performed in the rolling direction (RD) and in the transverse
direction (TD),
according to the standard EN ISO 7438:2005.
The bending ability was characterized by the ratio Alt between the radius of
curvature R of the bent sheet (in mm) and the thickness t of the sample (in
mm).
The results as summarized in Table 6 below.
Sample R/t (RD) R/t (TD)
(mm)
11 2.4 0.8 0.3
11 1.4 0.4 0.3
R1 2.2 2.7 2.7
R2 2.8 2.1 1.4
Table 6
In this table, t designates the thickness of the sample, and R/t designates
the
measured ratio between the radius of curvature of the bent sheet and the
thickness.
These results demonstrate that the steel according to the invention has an
improved
bending ability as compared to steels R1 and R2.
The Charpy energy of steels 11 and R2 was further determined on samples
collected
from the hot-rolled sheets, at temperatures ranging from -80 C to 20 C.
Especially, sub-size Charpy impact specimen (10 mm x55 mm x thickness of the
sheet) with V notches 2mm deep, with an angle of 45 and 0.25 mm root radius
were
collected from hot-rolled steel sheets made of steels 11 and R2.
At each temperature, the surface density Kcv of impact energy was measured. At
each temperature, the test was performed on two samples, and the average value
of the
two tests calculated.
The results are illustrated on Figure 15, and reported in Table 7 below.
In this table, T designates the temperature in degrees Celsius and Kcv
designates
the surface density of impact energy in J/cm2. In addition, the fracture mode
(ductile
fracture, mixed mode of ductile and brittle fracture or brittle fracture) is
reported.
As shown in Table 7 and Figure 15, the Charpy energy of steel 11 of the
invention is
much higher than the Charpy energy of steel R2. Moreover, the transition
temperature
from ductile to mixed fracture mode for steel 11 is lowered as compared to
steel R2.
Especially, in the steel of the invention, the fracture remains 100% ductile
at -20 C.
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28
Steel 11 Steel R2
T (thickness=1.45 mm) (thickness=1.8 mm)
( C) Kcv Fracture Kcv Fracture
(J/cm2) mode (J/cm2) mode
-80 33 mixed 3 brittle
-60 33 mixed 6 brittle
-40 35 mixed 15 mixed
-20 38 ductile 25 mixed
0 39 ductile 29 mixed
20 41 ductile 33 ductile
Table 7
These tests therefore demonstrate that the steel of the invention as an
improved
formability, ductility and toughness as compared to:
- steel R1, which has a Ti* content higher than 0.95% but was not cast under
the
form a thin product, and thus having TiC and coarse TiB2 precipitates,
- steel R2, which was cast in the form of a thin product but has a Ti* content
lower
than 0.95%, and thus having TiC and comprising may TiB2 precipitates with a
surface
area higher than 8 m2.
Finally, the mechanical properties of steels sheets 11, R1 and R2 were
determined.
Table 8 below reports the yield strength YS, the tensile strength TS, the
uniform
elongation UE, the total elongation TE and the elasticity modulus in tension
E, the work
hardening coefficient n and the Lankford coefficient r. Table 8 also reports
the volumic
percentage of TIB2 (fT,B2) precipitates for each steel.
YS TS UE TE fT1132 E
Sample n r
(MPa) (MPa) (%) (oh) (`)/0) (G Pa)
11 300 653 15.4 23.3 0.214 0.7 9 232
R1 245 530 14.2 19.7 0.192 0.8 12 240
R2 291 567 15.2 20.8 0.2 0.7 10.9 240
Table 8
These results demonstrate that the mechanical properties of steel 11 are
improved
as compared to the mechanical properties of steels R1 and R2. This improvement
is in
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29
particular due to the high proportion of very small size precipitates in steel
11, as
compared to steels R1 and R2.
The invention therefore provides a steel sheet and a manufacturing method
thereof
having at the same time a high elasticity modulus in tension, a low density,
and improved
castability and formability. The steel sheet of the invention can therefore be
sued to
produce parts with complex shapes, without inducing damages or surface
defects.
